Ceramic materials for gas separation and oxygen storage

ABSTRACT

A manganese oxide contains M1, optionally M2, Mn and O. M1 is selected from the group consisting of In, Sc, Y, Dy, Ho, Er, Tm, Yb and Lu. M2 is different from M1, and M2 is selected from the group consisting of Bi, In, Sc, Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb and Lu. These ceramic materials are hexagonal in structure, and provide superior materials for gas separation and oxygen storage.

CROSS-REFERENCE TO RELATED APPLICATION

The present application claims the benefit of U.S. ProvisionalApplication 61/407,580, filed 28 Oct. 2010, the entire contents of whichare hereby incorporated by reference, except where inconsistent with thepresent application.

FEDERALLY SPONSORED RESEARCH OR DEVELOPMENT

This invention was made with Government support under grant no.DMR-0706610 awarded by the National Science Foundation. The UnitedStates Government has certain rights in this invention.

BACKGROUND

The present invention relates to the selective storage and release ofoxygen and gas separation by ceramic materials. In particular, thepresent invention relates to methods of elevated temperature airseparation, oxygen storage, or any process related to temperature oroxygen partial-pressure dependent absorption and desorption of oxygenwith ceramic materials.

Recently ceramic materials have been increasingly researched due totheir reversible oxygen storage/release capacities (OSC) atelevated-temperatures. New ceramic materials for elevated-temperatureair separation are strong candidates to compete with cryogenicdistillation for commercial air separation and are also being researchedfor components to improve automotive exhaust catalysts, solar watersplitting, hydrogen-oxygen fuel cells, various non-aerobic oxidationprocesses, and assorted high-temperature production processes thatrequire high-purity oxygen (e.g. steel, copper, plastics, glass, etc.).Elevated temperature air separation methods have been projected to have20-30% less capital and operation cost, while being significantly moreenergy efficient, than conventional air separation methods. Thedevelopment of improved oxygen storage or carrier materials is alsocritical to the success of new energy related technologies such as“oxy-fuel” and “chemical looping” combustion systems for “clean coal”energy production, automotive pollution reduction, hydrogen-oxygen fuelcells, solar water splitting, and to improve the efficiency and cost ofvarious production processes (e.g. steel, copper, plastics), and theproduction of synthesis gas (H₂, CO) by partial oxidation of methane.

Ideal materials have large values of OSC (typically measured in moles ofoxygen per weight of material) and their absorption/desorption of oxygenoccurs over a narrow temperature range at near atmospheric conditions.Additional properties, such as oxygen partial pressure dependence ofabsorption/desorption, exothermic absorption and endothermic reduction,stability/recoverability in strong reducing conditions (e.g. CO and H₂atmospheres at high-temperatures), are also desired and being researchedfor various applications. Commercially, fluorite Ce_(1-x)Zr_(x)O₂compositions have been the recent ceramic OSC materials of choice forair separation, which function around 500° C. and have OSCs of ˜400-500μmol-O/g in oxygen atmospheres or as high as 1500 μmol-O/g with 20% H₂reversible reduction. Recent studies with Ce_(1-x)Cr_(x)O₂ have furtherboosted the OSC of the fluorite structure to as high as 2500 μmol-O/g inair and hydrogen atmospheres but require considerably higher reductiontemperatures (550-700° C.) and contain poisonous Cr⁶⁺. Currently,RBaCO₄O_(7+δ) (R═Y, Dy, Ho, Er, Tm, Yb, and Lu) andYBaCO_(4-x)Al_(x)O_(7+δ) have the best reported OSC at low-temperature,which have storage up to ˜2700 μmol-O/g and completely desorb at˜400-425° C. in O₂. The ease of reversible phase transitions between thehexagonal P6₃mc YBaCO₄O₇ and orthorhombic Pbc2₁ YBaCO₄O_(8.1) phases(which is a mixture of tetrahedrally and octahedrally coordinatedcobalt) is responsible for its oxygen storage behavior.

RMnO₃ (R=rare earths) and their competing hexagonal and perovskitecrystal structures have been studied for over fifty years.Conventionally, the formation of the perovskite phase versus thehexagonal phase is governed primarily by the size of the rare-earth ionin RMnO₃ (with constant Mn³⁺ size). During high-temperature solid statesynthesis in air, the perovskite phase forms easily with largerrare-earth elements (e.g. La, Pr, Nd, Sm, Gd, Tb, and Dy), while smallersize rare-earths (e.g. Ho, Er, Tm, Yb, Lu, and Y) favor the hexagonalphase. It has been observed that the perovskite phase is stable for atolerance factor,

${t = \frac{\left( {{R—}O} \right)}{\sqrt{2}\left( {{Mn}{—O}} \right)}},$

in the range of 0.855≦t≦1 (calculated at room temperature usingShannon's ionic size values), where the perovskite structure isincreasingly distorted as it approaches this lower limit and results inthe transition to the hexagonal phase at t<0.855. Recently, Zhou et al.suggested that the relative large difference in density between theperovskite and hexagonal phases can have a large impact on the formationof the perovskite versus the hexagonal near the lower limit of thetolerance factor. Regardless, DyMnO₃ and YMnO₃ have tolerance factors of0.857 and 0.854, respectively, and will tend to form the perovskite andhexagonal phases, respectively, under normal solid state reactionsynthesis. Thus the average (R—O) bond length of substituted samplescauses Dy_(1-x)Y_(x)MnO₃ to be on the cusp of this phase transition and,as further discussed herein, results in a mixed state under synthesis inair.

U.S. Patent Application Publication No. 2009/0206297 to Karppinen, etal. discloses an oxygen excess type metal oxide expressed with thefollowing formula (1) and exhibiting high speed reversible oxygendiffusibility whereby a large amount of excess oxygen is diffused at ahigh speed and reversibly in a low temperature region:

A_(j)B_(k)C_(m)D_(n)O_(7+δ)  (1)

where

A: one or more trivalent rare earth ions and Ca

B: one or more alkaline earth metals

C, D: one or more oxygen tetra-coordinated cations among which at leastone is a transition metal, where j>0, k>0, and, independently, m≧0, n≧0,and j+k+m+n=6, and 0<δ≦1.5. The metal oxide has high oxygendiffusibility and large oxygen non-stoichiometry at a low temperatureregion (500° C. or less, in particular 400° C. or less) and a ceramic isdisclosed for oxygen storage and/or an oxygen selective membranecomprised of the metal oxide. The Karppinen, et al. metal oxide has ahigh 2:1 ratio of expensive and poisonous Co (where C═Co) to the lessexpensive trivalent rare earth ions and alkaline earth metals. Inaddition, the compounds disclosed contain B═Ba that is highly reactivewith CO₂ and water vapor present in air and easily decompose when heatedjust above their optimal OSC temperature. The Karppinen, et al. metaloxide has the disadvantages of being expensive and not thermodynamicallystable or safe.

One disadvantage of currently used materials for oxygen storage or airseparation is that they depend on the creation of oxygen ion vacanciesor interstitial sites at high-temperatures in order to store the oxygen.Currently, the majority of materials for air separation usehigh-pressure (zeolites) or low-temperature (cryogenic distillation)methods consuming large amounts of energy. However, roughly over 80% ofcommercially produced oxygen is used in high-temperature industrialproductions process and many developing applications of OSC materialsoperate at high-temperatures as well. For any of these current andpotential systems, the redirection of the large amounts of waste heatgenerated from all these methods to ceramic OSC materials for onsite airseparation, would undoubtedly have potential net energy, economic, andwaste advantages versus conventional methods. The majority of newceramic OSC materials for such applications rely on the creation ofoxygen ion vacancies or interstitial sites at high-temperatures;however, this is a poor mechanism for currently known materials due tothe high-temperatures (up to 1000° C.) and large temperature gradients(˜300-800° C.) required for moderate oxygen storage capacities (lessthan 500 μmol-O/g).

Therefore, there remains a need for a method for the selective storageand release of oxygen that does not require extreme temperatures orlarge temperature gradients so as to reduce cost and energy consumption,as well as a method that is able to increase the amount of oxygen thatis able to be stored.

SUMMARY

In a first aspect, the present invention is a rare-earth manganeseoxide, comprising M1, optionally M2, Mn and O. M1 is selected from thegroup consisting of In, Sc, Y, Dy, Ho, Er, Tm, Yb and Lu. M2 isdifferent from M1, and M2 is selected from the group consisting of Bi,In, Sc, Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb and Lu. Mnand O are present in an atomic ratio of 1:z, and z is at least 3.15.

In a second aspect, the present invention is a rare-earth manganeseoxide, comprising M1, M2, Mn and O. M1 is selected from the groupconsisting of In, Sc, Y, Dy, Ho, Er, Tm, Yb and Lu. M2 is different fromM1, and M2 is selected from the group consisting of Bi, In, Sc, Y, La,Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb and Lu. M1 and M2 arepresent in an atomic ratio of x:1−x, and x=0.1 to 0.9.

In a third aspect, the present invention is a rare-earth manganeseoxide, comprising (i) Mn, having a formal oxidation state between 2 and3, or between 3 and 4, (ii) O, and (iii) at least one element selectedfrom the group consisting of Bi, In, Sc, Y, La, Ce, Pr, Nd, Sm, Eu, Gd,Tb, Dy, Ho, Er, Tm, Yb and Lu. The rare-earth manganese oxide has anaverage temperature of maximum oxygen absorption upon heating andcooling, T_(maxA), of at most 400° C., and a temperature of maximumoxygen desorption, T_(maxD), of at most 400° C.

In a fourth aspect, the present invention is an oxygen conductingmembrane, comprising (1) a rare-earth manganese oxide, and (2) a supportmaterial. The membrane has first and second opposing surfaces, themembrane is not permeable to nitrogen gas, the rare-earth manganeseoxide forms a contiguous structure exposed on both the first and secondopposing surfaces, and the rare-earth manganese oxide has an averagetemperature of maximum oxygen absorption upon heating and cooling,T_(maxA), of at most 400° C., and a temperature of maximum oxygendesorption, T_(maxD), of at most 400° C.

In a fifth aspect, the present invention is an oxygen conductingmembrane, comprising (1) a manganese oxide, and (2) a support material.The membrane has first and second opposing surfaces, and the membrane isnot permeable to nitrogen gas. The manganese oxide forms a contiguousstructure exposed on both the first and second opposing surfaces, andthe support material comprises at least one member selected from thegroup consisting of an organic polymer, a silicone rubber and glass.

In a sixth aspect, the present invention is a method of preparingoxygen, comprising separating oxygen from a mixture of gases containingthe oxygen, by conducting the oxygen through the manganese oxide, orabsorbing and releasing the oxygen from the manganese oxide.

In a seventh aspect, the present invention is a method of catalyzing areaction with oxygen, comprising catalyzing the reaction with themanganese oxide.

In an eighth aspect, the present invention is a method generatingelectricity, comprising burning a carbon-containing fuel with oxygen, ina generator or power plant, wherein the oxygen is prepared by using themanganese oxide.

DEFINITIONS

The terms “rare-earth”, “rare-earth element”, and “rare-earth metal”include Bi, In, Sc, Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm,Yb and Lu.

The formal oxidation state of manganese, Mn, in a manganese oxide, forexample a rare-earth manganese oxide, may be determined by: (a)multiplying the relative amount of each element other than manganese andoxygen by the most common oxidation of that element in an oxide, (b)multiplying the relative amount of oxygen by 2, (c) subtracting thefirst value (a) from the second value (b), and then dividing by therelative amount of manganese. The most common oxidation state in anoxide of Bi, In, Sc, Y, La, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yband Lu is 3; and the most common oxidation state in an oxide of Ce andTh is 4. In the case of transition metals which have two or three commonoxidation states in an oxide, the most common oxidation state in anoxide is the largest common oxidation state in an oxide; example valuesinclude 2 for Cu and Ni, 3 for Co and Fe, 4 for Ti, 5 for V and Nb, 6for Mo and W, and 7 for Re. The following is an exemplary calculationfor Dy_(0.3)Y_(0.7)MnO_(3.25): a=[(3×0.3)+(3×0.7)]=3; b=2×3.25=6.5;c=6.5−3=3.5; formal oxidation state of Mn=3.5/1=3.5.

The average temperature of maximum oxygen absorption upon heating andcooling, T_(maxA), is the temperature where the first derivative ofoxygen content as a function of temperature is a maximum, as measured bythermogravimetric analysis in pure O₂ with heating and cooling rates of0.1° C./minute. The manganese oxides may have an average temperature ofmaximum oxygen absorption upon heating and cooling, T_(maxA), of at most400° C., of at most 300° C., or of at most 250° C.

The temperature of maximum oxygen desorption, T_(maxD), is thetemperature where the first derivative of oxygen content as a functionof temperature is a minimum, as measured by thermogravimetric analysisin pure O₂ with heating and cooling rates of 0.1° C./minute. Themanganese oxides may have a temperature of maximum oxygen desorption,T_(maxD), of at most 400° C., of at most 300° C., or of at most 250° C.

BRIEF DESCRIPTION OF THE DRAWINGS

FIGS. 1A-1D are X-ray diffraction patterns of the ceramic material;

FIGS. 2A-2E are graphs of oxygen absorption/desorption of the ceramicmaterial;

FIG. 3 is a graph of oxygen content versus temperature with TGAannealing and reduction of the ceramic material;

FIGS. 4A and 4B are graphs of TGA oxygen content versus temperature forthe ceramic material with heating (4A) and cooling (4B);

FIG. 5 is a graph showing relevant temperatures for the ceramicmaterial;

FIG. 6 is a graph of TGA reduction in 21% O₂ of high-pressure annealedceramic material to stable 3.0 oxygen content;

FIG. 7 is a graph of TGA of the ceramic material switching between Arand O₂ at various isotherms;

FIG. 8 is a graph showing phase mapping of the ceramic material;

FIGS. 9A and 9B are graphs of TGA (9A) and dilatometry (9B) measurementsfor DyMnO₃ in 21% O₂;

FIGS. 10A and 10B are graphs of (10A) TEC values for P6₃ cm and Hex₂phases and (10B) chemical expansion (CE) parameter during the Hex₂-P6₃cm phase transition for Dy_(1-x)Y_(x)MnO_(3+δ);

FIG. 11 is a dilatometry measurement of perovskite DyMnO₃ in 21% O₂;

FIGS. 12A-12F are XRD patterns of DyMnO_(3+δ) with δ=−0.037, 0.0, 0.18,0.21, 0.24, and 0.35, arrows indicate the increase and decrease of peakintensity for hexagonal phases;

FIG. 13 is a graph of TGA reduction in 21% O₂ of high-pressure annealedDyMnO_(3+δ) to stable 3.0 oxygen content;

FIG. 14 is a graph of XRD comparison of DyMnO_(3+δ) samples: P6₃ cm(δ=0), nearly single phase Hex₂ (δ=0.24), and mixed phase of Hex₂ andHex₃ (δ=0.35);

FIG. 15 is a graph of NPD pattern for DyMnO_(2.963) (SEPD), plus signsare observed data and the line below is the difference betweenexperimental data and best fit calculated from the Rietveld refinementmethod;

FIGS. 16A and 16B are best-fit Rietveld refinement patterns using highresolution synchrotron X-ray data with a wavelength of 0.40225 Å (11BM-B), observed (plus signs) and calculated (solid line) intensities aredisplayed together with their difference (solid line at the bottom ofeach panel), lower and upper tick marks indicate the locations of Braggreflections for the parent P6₃ cm and superstructure R3 phases,respectively; and

FIGS. 17A-17D show TGA annealing: (17A) in air of DyMnO_(3+δ) afterinitial synthesis in argon and subsequent reduction in H₂ to Dy₂O₃ andMnO; (17B) in oxygen of YMnO_(3+δ) after initial synthesis in air andreduction in H₂ after anneal at 190 atm. of oxygen; (17C) in air forperovskite La_(0.5)Sr_(0.5)Fe_(0.5)CO_(0.5)O_(3+δ); and (17D) in O₂ forperovskite LaMnO_(3+δ).

FIG. 18 are graphs of TGA oxygen content versus temperature for theceramic materials DyMnO_(y), YMnO_(y), and HoMnO_(y) with heating andcooling in oxygen.

DETAILED DESCRIPTION

The present invention provides a new system of ceramic materials forelevated temperature air separation methods, oxygen storage, or anyprocess related to temperature or oxygen partial-pressure dependentabsorption and desorption of oxygen. These processes are, but notlimited to, Thermal Swing Absorption (TSA) and Ceramic AutothermalRecovery (CAR) methods.

The materials of the present invention behave like an “oxygen sponge.”Just as a sponge can absorb and release water under different pressures,the materials of the present invention can absorb and release oxygenwhen exposed to different temperatures or gasses. This property can beused to separate the major components of the air in the atmosphere,oxygen and nitrogen, by storing the oxygen in the ceramic materials,leaving nitrogen in the atmosphere. The oxygen that is absorbed andstored in the material is preferably oxygen ions, O²⁻. The ceramicmaterials are able to selectively absorb and release oxygen with near100% selectivity and not absorb other gases. This is a key property thatmakes the materials excellent OSC materials.

The present invention provides for a ceramic material system that isrepresented by the formula

A_(j)B_(k)C_(m)D_(n)O_(3+δ)

where

A: one or more trivalent rare earth ions and tetravalent rare earthelements,

B: one or more alkaline earth metals and Pb, Bi, In and Sc,

C, D: one or more oxygen bi-pyramidally-coordinated cations among whichat least one is a transition metal or post-transition metal, where j>0,k≧0, and, independently, m≧0, n≧0, and j+k=1, m+n=1, and 0<δ≦0.5. Inother words, C and D can be any transition metal or post-transitionmetal, for example, groups 3 through 12 on the periodic table (Sc, Ti,V, Cr, Mn, Fe, Co, Ni, Cu, Zn, Y, Zr, Nb, Mo, Tc, Ru, Rh, Pd, Ag, Cd,Hf, Ta, W, Re, Os, Ir, Pt, Au, Hg, Rf, Db, Sg, Bh, Hs, and Cn) and Ga,In, Sn, Tl and Pb. The metal oxide has high oxygen diffusibility andlarge oxygen nonstoichiometry at a low temperature region of 400° C. orless and a ceramic is disclosed for oxygen storage and/or an oxygenselective membrane comprised of the metal oxide.

A and B can be chosen from the following ions: 3+ ions: Y, La, Pr, Nd,Sm, Eu, Gd, Tb, Ho, Dy, Er, Tm, Yb, Lu, Bi, In and Sc; 4+ ions: Th andCe; and 2+ ions: Ca, Sr, Ba and Pb.

Alternatively, the ceramic materials are manganese oxides, for examplerare-earth manganese oxide, containing M1, optionally M2, Mn and O. M1is selected from Sc, Y, Dy, Ho, Er, Tm, Yb and Lu; M2 is different fromM1, and M2 is selected from Sc, Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy,Ho, Er, Tm, Yb and Lu. In some compositions, Mn and O are present in anatomic ratio of 1:z, and z is at least 3.1. In some compositions M1 andM2 are present in an atomic ratio of x:1−x, and x=0 to 1. In somecompositions Mn has a formal oxidation state between 3 and 4.Preferably, the manganese oxide has an average temperature of maximumoxygen absorption upon heating and cooling, T_(maxA), of at most 400°C., and a temperature of maximum oxygen desorption, T_(maxD), of at most400° C.

It is also possible to substitute other metals for Mn in the ceramicmanganese oxides. For example, 10 atomic %, or 15 atomic % of the Mncould be substituted with Co, Ni, Fe, Cu, Ru, Rh and/or In.

The value of z, which corresponds to the atomic ratio of oxygen permanganese, may be at least 3.2, or at least 3.24, or 3.1 to 3.4. It isalso possible to remove oxygen, and have a value of z which is less than3, for example where z is at most 2.9, or at most 2.8, such as 2.88 to2.80. Preferably, the manganese oxide has an average temperature ofmaximum oxygen absorption upon heating and cooling, T_(maxA), of at most400° C. (including at most 300° C., and at most 250° C.), and atemperature of maximum oxygen desorption, T_(maxD), of at most 400° C.(including at most 300° C., and at most 250° C.).

The hexagonal DyMnO₃ and YMnO₃, (materials with no excess oxygen (δ=0))are previously known compounds with known crystal structures. Theceramic material system is a new material having a new crystal structurewith excess oxygen (as indicated by δ>0, these excess oxygen ions belongto the crystal structure. Their inclusion in crystal structure as O²⁻ions is compensated by oxygenation of Mn ions from 3+ to 3+2δ, such thatthe overall charge neutrality is preserved). In the example below,Dy_(1-x)Y_(x)MnO_(3+δ) is synthesized, wherein based on the aboveformula, A=Dy, B═Y, and C and D=Mn, and 0≦x≦1.

The present invention provides for a method of making the ceramicmaterials by synthesizing hexagonal P6₃ cm material, and oxygenating thematerial in partial-pressures of oxygen at low elevated temperatures(200-300° C.). This is also further described in the example below.

Unlike other oxygen storage materials, which depend on the creation ofoxygen ion vacancies at high-temperatures, this system relies on areversible phase transition between δ=0 and δ=0.25-0.5 phases at lowertemperatures of approximately 300° C. They exhibit large changes ofoxygen content over both a narrow temperature range and a smalldifference of oxygen pressure near atmospheric conditions. Theseattributes of this system allow use of inexpensive processes toincorporate and extract large quantities of oxygen.

An oxygen conducting membrane may be prepared from the manganese oxides,in combination with a support material. Because the oxygen absorptionand conduction occur at temperatures much lower than in the perovskitematerials, a much larger variety of support materials may be used.Possible support materials include materials which decompose whenexposed to air at a temperature of 500° C., 400° C. or even 300° C., orwhich have a glass transition temperature or a melting point of at most500° C., at most 400° C. or at most 300° C. Specific examples includeorganic polymers, silicone rubbers, glass, graphite, carbon black,aluminum, copper, iron, nickel, steel, zinc, tin, lead and alloysthereof. The manganese oxide forms a contiguous structure exposed onboth opposing surfaces of the membrane. It may also be desirable for thesupport material to form a contiguous structure exposed on both opposingsurfaces of the membrane, especially in the case of an electricallyconductive support (see, for example, Thorogood et al., U.S. Pat. No.5,240,480).

The present invention provides for a method of storing oxygen, includingthe steps of exposing the ceramic system to oxygen, containing gas,selectively absorbing the oxygen in the system, and storing the oxygen.It has been discovered that these hexagonal materials have unusuallylarge oxygen absorption at approximately 200-300° C. in oxygenatmospheres (storage, excess oxygen preserved on cooling to roomtemperature). The following steps are involved in the process of oxygenstorage: oxygen molecules 2 present in a gas (for example air) reach thesurface of the material where they are split to oxygen ions, the oxygenions then diffuse through the crystal lattice of material and congregatenear ions of the system such as manganese to form newly discoveredcrystal structures (δ>0).

These drastic uptakes of oxygen were observed to completely desorb whenmaterials transitioned back to the stoichiometric P6₃ cm state (δ=0)during increased heating of the system to 275-375° C. or changing tolower oxygen partial-pressures (release) surrounding the system. Thesteps involved in the process of oxygen release occur in reverse order:oxygen ions diffuse towards the material surface, oxygen ions recombineon the surface to form molecular O₂, which can be extracted and used.Therefore, the present invention also provides for a method of releasingoxygen, including the step of releasing the oxygen absorbed in theceramic system.

For example, these materials can be used with known processes for airseparation with OSC materials such as TSA and CAR. Thermal SwingAbsorption (TSA) relies on temperature dependent oxygenabsorption/desorption of its “oxygen carrier”. In this method, multiplebeds of sorbent cycle in between two chambers that are at differenttemperatures. This creates oxygen rich and oxygen deficient atmospheresin each chamber. More recently, a method was patented in 2000 by Lin etal. (U.S. Pat. No. 6,059,858) for perovskite materials, which combinesTSA and PSA (Pressure Swing Absorption) techniques in a process namedCeramic Autothermal Recovery (CAR). Again, multiple beds filled withsorbent are cycled through two chambers with a temperature gradient;however, in this method the chambers are also at different oxygenpartial-pressures. Again, this creates two chambers that are oxygen richand deficient. Sorbents designed for CAR also have endothermic reductionand exothermic absorption; therefore, the process operatesautothermally, needing little or no heat added once operational.

The present invention provides for a method of separating gaseous ormolecular O₂ from at least a second gas, by exposing the ceramic systemto the gaseous or molecular O₂ and second gas, absorbing the gaseous ormolecular O₂ into the system, and separating the gaseous or molecular O₂from the second gas. This method can be performed with any number ofgasses present, and allows for separation of gaseous or molecular O₂from the other gasses. Preferably, the second gas is nitrogen inseparating oxygen from air. Other second gasses include hydrogen, He,Ne, Ar, Kr, Xe, Rn, N₂, CO, CO₂, CH₄.

In the example below, polycrystalline samples ofDy_((1-x))Y_((x))MnO_(3+δ) were synthesized by solid state reaction withappropriate amounts of Dy₂O₃, Y₂O₃ and MnO₂ (all with >99.99% purity).For all samples, reactants were thoroughly mixed in an agate mortar, andfired in air in the temperature range of 800-1200° C. with intermediategrindings followed by pressing samples into high-density pellets. Allsteps of the synthesis were monitored with X-ray powder diffractionmeasurements. Samples were fired several times until single phaseperovskite was obtained, except for YMnO₃, which forms hexagonalstructure under these conditions (FIG. 1: X-ray diffraction patterns forDyMnO₃ (a) and YMnO₃ (b)). Similar to YMnO₃ the HoMnO₃ and ErMnO₃ formhexagonal phases in air. Hexagonal phases were acquired from perovskitesamples by firing under ultra-high-purity Argon (99.999%) with ahydroxyl purifier (measured oxygen partial-pressures of 5 to 10 ppm) ina temperature range of 1200-1400° C. (FIG. 1: X-ray diffraction patternfor DyMnO₃ (c)). Hexagonal samples were then oxygenated at ambientpressure or under 250-350 bars of oxygen pressure at 200-400° C.followed by cooling at 0.1°/min to room temperature in order to achievethe largest oxygen content possible (FIG. 1: X-ray diffraction patternfor DyMnO_(3.35) (d)).

Temperature and oxygen partial-pressure dependence of reversible oxygenstorage capacities (OSC) were demonstrated by thermogravimetric analysis(FIG. 2 reversible oxygen absorption/desorption of a typical perovskitematerial considered for application Sr₂FeCoO_(5-δ) (a, as evidenced bythe data below in the Example), reversible oxygen absorption/desorptionof the hexagonal material (b, as function of temperature) and (c, as afunction of oxygen pressure), best available oxygenabsorption/desorption material from literature: T. Motohashi, S. Kodota,Mater. Sci. Eng. B 148 (2008) 196 (d) and (e)). Hexagonal manganiteshave been largely believed to remain stoichiometric in oxygen content atelevated-temperatures; however, the thermogravimetric measurements ofoxygen annealed hexagonal samples indicated unusually large oxygenabsorption over a narrow temperature range ˜200-300° C., which return tostoichiometric behavior above 275-375° C. in O₂ atmosphere. In additionto temperature dependence, the oxygen content of Dy_(1-x)Y_(x)MnO_(3+δ)was also found to be sensitive to changes in partial pressures of oxygenin these temperature ranges. The hexagonal P6₃ cm phase of this systemwas found to have considerable stability at high-temperature in partialpressures of oxygen and to be recoverable from a reduced state withnegative values of δ obtained from reduction in hydrogen at 400° C.

The ceramic system of the present invention can be used for elevatedtemperature gas separation and oxygen storage methods, which include,but are not limited to, oxygen and nitrogen production and componentsfor oxy-fuel “clean coal” power plants, automotive exhaust catalysts,H₂—O₂ fuel cells, solar water splitting methods, and steel, copper, andplastic production and any other various industrial production processeswhich require high-purity oxygen and have large amounts of waste heat.The ceramic system can replace cryogenic distillation or pressure swingabsorption for commercial air separation.

The ceramic system of the present invention has several advantages overthe prior art. This system of materials has been shown to havecomparable OSC with current commercial ceramic materials while operatingat lower temperatures and has a smaller necessary temperature gradientfor oxygen absorption/desorption. This system can also have fasteroxygen absorption/desorption rates. In addition it is made ofinexpensive and abundant elements.

The invention is further described in detail by reference to thefollowing experimental examples. These examples are provided for thepurpose of illustration only, and are not intended to be limiting unlessotherwise specified. Thus, the invention should in no way be construedas being limited to the following examples, but rather, should beconstrued to encompass any and all variations which become evident as aresult of the teaching provided herein.

EXAMPLES Example 1 Experimental Techniques

Synthesis was done by solid state reaction, which is further detailed inthe following section. X-ray powder diffraction (XRD) measurements weremade with a Rigaku D/MAX powder diffractometer in the 2θ=20-70° rangewith CuKα radiation. Thermogravimetric analysis (TGA) measurements weremade with Cahn TG171 and Cahn TherMax700 thermobalances in severaldifferent partial-pressures of oxygen and hydrogen (balanced with argon)up to 1400° C. at heating and cooling rates of 0.1-1.0°/min. TGA sampleswere approximately 1 g and were measured with a 5 μg precision.Dilatometry measurements were made with a Linseis DifferentialDilatometer L75 and samples were measured with a 1 μm precision.

Results and Discussion

Synthesis and Stability

Polycrystalline samples of hexagonal Dy_(1-x)Y_(x)MnO_(3+δ) weresynthesized by solid state reaction with appropriate amounts of Dy₂O₃,Y₂O₃, and MnO₂ (all with >99.99% purity). For all samples, reactantswere thoroughly mixed in an agate mortar, and fired in air in thetemperature range of 800-1300° C. with intermediate grindings followedby pressing samples into high-density pellets at approximately 1 kbar.All steps of the synthesis were monitored with XRD measurements andcompared to previous diffraction measurements in the literature ofhexagonal P6₃ cm and perovskite Pnma phases of DyMnO₃ and YMnO₃ (FIG.1). Dy_(1-x)Y_(x)MnO_(3+δ) samples which formed the perovskite or amixed phase in air instead of the single phase hexagonal structure (x=0,0.1 0.3, 0.5, 0.7), were then fired under ultra-high-purity argon(99.999%) at 1300 and 1400° C. Dy_(1-x)Y_(x)MnO_(3+δ) samples (x=0 and0.1) were then subsequently fired under ultra-high-purity argon with ahydroxyl purifier (oxygen partial pressures of 5-10 ppm) at 1400° C. Allsamples achieved the hexagonal P6₃ cm structure after these conditions.

Considerable effort was devoted to synthesizing Dy-rich, homogenoushexagonal samples. The hexagonal DyMnO₃ phase has been previouslyachieved by epitaxially stabilized crystal growth with thin-films,thermal decomposition with polynuclear coordination compound precursors,quenching methods from 1600° C. in air or 1250° C. in argon for 3 dayswith sol-gel methods. This work confirmed that synthesis in argon athigh-temperature tends to favor the formation of the hexagonal phase,while synthesis in oxygen tends to favor the perovskite phase.

FIG. 8 is a mapping of the phases that were measured with XRD afterseveral synthesis steps, which clearly shows that increasing reducingconditions are needed to form the hexagonal phase as the average ionicradius of the R-site increases. The oxygen content dependence of thetolerance factor, which was previously studied for substituted SrMnO₃,is most likely responsible for this behavior. The formation of oxygenvacancies in RMnO_(3+δ) (δ<0) causes a change in oxidation state in someof the Mn³⁺ cations to Mn²⁺, resulting in a net Mn^((3+2δ)+) cation,which increases the (Mn—O) bond length with decreasing δ. TGAmeasurements in oxygen (FIGS. 4A and 4B) show the reduced oxygencontents after synthesis of single phase hexagonal samples in argon. Theresulting larger (Mn—O) bond lengths of these samples decrease theirtolerance factor below the lower limit of 0.855 and results in theperovskite phase undergoing a phase transition to the hexagonal phase.Using Shannon room temperature values, the minimum necessary value of δranges from −0.023 to −0.0027 to have t≦0.855. It was observed, however,that samples with the corresponding δ values did not transformcompletely to the hexagonal phase (TABLE 1). Previous in situmeasurements with Ca and La substituted SrMnO₃ have shown that both (Ca,Sr, La—O) and (Mn—O) bond lengths increase with temperature in a mannerwhich increases the value of the tolerance factor. Therefore, thetransition from the perovskite phase to the hexagonal phase most likelyoccurs in various oxygen pressures at 6 which is a function oftemperature, which occurs for DyMnO_(3+δ), for example, in ˜10 ppm O₂ at1400° C. as observed here or in air at 1600° C. as previously reported.The combination of previous in situ measurements with similar manganitesand XRD measurements of various oxygen contents after progressiveincreased reducing conditions strongly support this conclusion.

TABLE 1 Hexagonal-Perovskite transition δ-values: from conditiont(δ_(Theo)) = 0.855, δ_(obs.) are observed values from TGA, and valuesof t are calculated with Shannon values. x δ_(Theo.) δ_(obs.)t(δ_(obs.)) 0 −0.0230 −0.037(0) 0.852(8) 0.1 −0.0201 −0.049(3) 0.847(0)0.3 −0.0143 −0.015(1) 0.854(9) 0.5 −0.0085 −0.020(9) 0.853(1) 0.7−0.0027 −0.017(1) 0.852(7)

Several other factors can also affect this transition. It can beenhanced by the difficulty of maintaining the twelvefold coordination ofR required for the perovskite in a high-temperature, oxygen deficientatmosphere; thus, an eightfold coordination with hexagonal symmetryresults. As mentioned in the introduction, the relative large differencein density between the perovskite and hexagonal phases plays asignificant role in this transition. The crystal strain of theperovskite phase by Jahn-Teller distortions may also destabilize thestructure to favor the hexagonal phase. In any case, the reducingconditions needed for production of bulk polycrystalline samples ofDyMnO₃ and Dy_(0.1)Y_(0.9)MnO₃ by standard firing methods were very nearto decomposition to simple oxides and many attempts were needed to findthe most favorable temperature and length of the firings. Increasedsubstitution of Y in DyMnO₃ considerably eases the necessary reducingconditions to synthesize the hexagonal phase.

The stability of hexagonal Dy_(1-x)Y_(x)MnO_(3+δ) compounds was alsotested by firing samples at high-temperatures, 1100-1400° C., in oxygen.As reducing conditions favor the hexagonal phase, atmospheres that allowsamples to remain near stoichiometric in oxygen content (or yield excessoxygen content, δ>0) at high-temperature promotes the perovskite phaseover the hexagonal, due to the smaller size of the Mn^((3+2δ)+) cationin oxygen versus argon. Dy rich samples (x=0, 0.1) began slightdecomposition back to the perovskite at 1100° C. and completelytransformed back to the perovskite at 1400° C. The remaining samples(x=0.7, 0.5, 0.3, 0.1, 0) remained hexagonal with no signs ofdecomposition back to the perovskite up to 1400° C. These results are inagreement with the presented tolerance factor arguments and can alsoexplain why small rare-earth manganites (R═Y, Ho, Er, Tm, Yb, and Lu)have been observed to transition to the perovskite phase underhigh-pressure oxygen, while smaller A-site cations (R═Sc and In) willnot transform to perovskite under similar conditions.

Thermogravimetric Measurements of OSC

After initial synthesis of the hexagonal phase, all samples wereannealed in TGA up to 500° C. with 0.1-1° C./min heating and cooling invarious partial-pressures of oxygen and hydrogen to measure OSC valuesand to demonstrate temperature and oxygen-partial pressure dependence ofoxygen content. The oxygen content after initial synthesis ofDyMnO_(3+δ) and YMnO_(3+δ) were determined with TGA by the difference inweight between oxygenated samples and their respective reductionproducts, Dy₂O₃, Y₂O₃, and MnO (verified by XRD), obtained by firstannealing at 1° C./minute in O₂ and followed by slow reduction at 0.1°C./minute in 42% H₂/Ar as shown for DyMnO_(3+δ) in FIG. 3. DyMnO₃ andYMnO₃ were observed to reduce to stable stoichiometric P6₃ cm phase inoxygen above 375 and 275° C., respectively. Using this information,stable weights of all samples above 400° C. in O₂ in TGA were normalizedto δ=0. TABLE 2 is a compilation of OSC values achieved by the followingassorted methods.

TABLE 2 OSC (μmol-O/g) of Dy_(1−x)Y_(x)MnO_(3+δ); isotherms in AR—O₂were done near “transition temp.” on FIG. 5 and reduction in H₂ wereconducted at 400° C. (*calculated values). Theoretical cooling rate 1.0°C./min 0.1° C./min 0.1° C./min Isotherm Max. Isotherm atmosphere O₂ O₂250 bars O₂ Ar—O₂ Ar—O₂/O₂ - with H₂ reduction x 0.0 812 926 1327 7701884* +455  0.1 729 1138 1388 — 1937* +499* 0.3 444 1200 1397 1149 2055*+596* 0.5 637 1169 1625 1091 2187* +705  0.7 133 849 1542 493 2337*+829* 0.9 176 952 1694 — 2501* +972* 1.0 53 666 1338 95 2606* +1051 

The TGA data on heating of DyMnO_(3+δ) in air (FIG. 17A) clearly showsthe reversible absorption (around 250-300° C.) and desorption (above320° C.) of excess oxygen in a narrow temperature range. On coolingabsorption occurs around 280° C. The resulting OSC values measured bythe difference in oxygen content between the stoichiometric phaseobserved above 400° C. (δ=0) and the oxygen content around 200° C.(δ=0.01-0.29) yielded values of 54-1200 μmol-O/g forDy_(1-x)Y_(x)MnO_(3+δ) compounds. Similar TGA measurements forYMnO_(3+δ) (FIG. 17B) in oxygen showed smaller amounts of oxygenabsorption (around 180-240° C.) and desorption (above 260° C.) occurringat lower temperatures. As a comparison, FIGS. 17C-17D show typical TGAtraces for perovskite materials La_(0.5)Sr_(0.5)Fe_(0.5)Co_(0.5)O_(3+δ)and LaMnO_(3+δ). FIG. 18 shows exemplary oxygen adsorption anddesorption on heating and reversible adsorption on cooling in oxygenatmosphere of hexagonal YMnO₃ and DyMnO₃. These materials show largeoxygen non-stoichiometry up to δ=−0.50, indicating possibly very largeOSC. However, reaching these values require a wide temperature gradientsexceeding 700° C., unlike the hexagonal compounds where δ=0.30-0.40 canbe achieved after high oxygen pressure anneal over narrow temperatureranges less than 100° C.

Samples (x=0.1, 0.3, 0.5) were able to attain the highest oxygencontents while having smaller molar weights resulting in OSC values upto 2000 μmol-O/g in air after prolonged annealings. XRD data of thesephases for Dy_(1-x)Y_(x)MnO_(3+δ) indicate the formation of superstructures at δ=0.25 (Hex₂) and δ=0.40 (Hex₃), which are currentlystudied with neutron powder diffraction measurements.Dy_(1-x)Y_(x)MnO_(3+δ) materials showed also oxygen content dependenceon oxygen partial pressure at constant temperatures. Cycling between O₂and Ar atmospheres for periods of 12 hours at temperatures as low as300° C. yielded OSC values of 95-1149 μmol-O/g. While the valuesreported here do not surpass the best observed OSC in the literature,the Dy_(1-x)Y_(x)MnO_(3+δ) system does have several advantages forapplication over other candidate materials. Dy_(1-x)Y_(x)MnO_(3+δ)system has the lowest reported absorption/desorption cyclingtemperatures, being approximately 100° C. lower than the lowest reportedtemperatures of YBaCO_(4-x)Al_(x)O_(3+δ) while showing far superiorthermodynamical stability. Additionally, from a hazardous waste and coststandpoints, mass-production of manganese oxides is much preferable tothat of cobalt or chromium oxides. Finally, unlike the majority of otherOSC materials, which depend on the creation of oxygen ion vacancies athigh temperatures, the hexagonal Dy_(1-x)Y_(x)MnO_(3+δ) (similar toYBaCO_(4-x)Al_(x)O_(7+δ)) relies on reversible phase changes due tooxygen filling/discharge of the interstitial sites at much lowertemperatures.

Temperature dependence of oxygen content of Dy_(1-x)Y_(x)MnO_(3+δ) wasmeasured in TGA with heating and cooling rates of 0.1 and 1.0° C./minuteunder high-purity oxygen. The resulting TG curves (0.1° C./min, FIGS. 4Aand 4B) clearly show the reversible absorption and desorption of oxygenbelow 400° C. in a narrow temperature range. OSC values were measured bythe difference in oxygen content between the stoichiometric P6₃ cm phaseobserved above 400° C. (δ=0) and the final oxygen content after cooling(δ=0.01-0.29), which yielded a large range of values, 54-1200 μmol-O/g(TABLE 2). Comparing 0.1 versus 1.0° C./min, resultant TGA curves andOSC values indicate that oxygen absorption rates increase with increasedDy content. Yet, samples (x=0.1, 0.3, 0.5) were able to achieve higheroxygen content than the pure Dy sample on 0.1° C./min cooling. Y-richsamples (x=0.7, 0.9, and 1) were also able to yield larger OSC valuesthan observed in TGA with long isothermal steps with slow cooling andindicate, if given enough time (>24 hours), would reach excesses inoxygen content up to δ=0.25. Four different temperatures were alsoidentified from TGA runs in O₂, which are plotted in FIG. 5: the averagetemperature of maximum oxygen absorption on heating and cooling

$\left( {\frac{d\left( {{Ox}.\; {Cont}.} \right)}{d\left( {{Temp}.} \right)} = {{local}\mspace{14mu} {maximum}}} \right)$

maximum oxygen desorption

$\left( {\frac{d\left( {{Ox}.\; {Cont}.} \right)}{d\left( {{Temp}.} \right)} = {{local}\mspace{14mu} {minimum}}} \right),$

transition temperature from oxygen absorption to desorption, and thetemperature where samples return to the stoichiometric P6₃ cm phase

$\left( {\frac{d\left( {{Ox}\; {{Cont}.}} \right)}{d\left( {{Temp}.} \right)} = 0} \right)$

(these can be approximately identified on FIGS. 4A and 4B byinspection). A thermal swing absorption process for air separation foreach of these samples can involve cycling in between their respectivetemperatures slightly above “Ox=3.0” and slightly below “Ave. Max.Absorption” in air. These resulting cycling ranges are approximately220-300° C. (x=1) to 310-390° C. (x=0) and produce large amounts of O₂over these narrow temperature ranges.

Samples were also annealed at 250 bars of O₂ at 400° C. followed by 0.1°C. cooling. The oxygen content of these samples after annealing weredetermined in TGA by the difference in weight between their startingweight and their weight at 375° C. (1° C./min heating) in 21% O₂normalized to δ=0 (FIG. 6). All samples showed significant increase inOSC (particularly with samples rich in Y content) under high pressureversus identical cooling in 1 bar of O₂ (TABLE 2). FIG. 6 also showsincreased stability of oxygen content on reduction at ˜300° C. for allsamples, which shows Dy_(1-x)Y_(x)Mn³⁺ _(0.5)Mn⁴⁺O_(3.25) (Hex₂) is astable phase and another stable phase at or above an oxygen content of3.35 may exist. XRD data of these phases for DyMnO_(3+δ) have beenpreviously reported and indicate the formation of super structures atδ=0.25 (Hex₂) and δ=0.40 (Hex₃). Though these samples show increasedoxygen content from atmospheric pressure oxygenations, the Mn³⁺ cationis still not completely oxidized to the Mn⁴⁺ state, which is ideal formaximum OSC values. TABLE 2 also includes these theoretical values ofOSC for a reversible Mn³⁺-Mn⁴⁺ (δ=0-δ=0.5) transition. The significantincrease of calculated OSC values with increased Y content in TABLE 2 isdue to the smaller molar weight of Y cation and is one of the reasonswhy the Dy_(1-x)Y_(x)MnO_(3+δ) system was chosen for study after workwith hexagonal

DyMnO_(3+δ).

Oxygen partial-pressure dependence of oxygen content ofDy_(1-x)Y_(x)MnO_(3+δ) and absorption/desorption reversibility weredemonstrated with TGA measurements at isotherm in cycling O₂ and Aratmospheres every ˜12 hours (FIG. 7). Samples were held at temperaturesnear their respective “transition temperatures” defined from FIG. 5 (forx=0, 0.3, 0.5, 0.7, and 1; T=330, 300, 280, 250, and 230° C.,respectively) and yielded OSC values of 95-1149 μmol-O/g (TABLE 2).Besides DyMnO_(3+δ), which clearly comes to equilibrium in O₂, these OSCvalues are comparisons of absorption after 12 hours. Given more time,these samples can achieve higher oxygen content, for example δ=0.28 wasobtained for Dy_(0.3)Y_(0.7)MnO_(3+δ) after ˜60 hours. Isothermalmeasurements also show oxygen content to have asymptotic behaviorsignificantly lower than achieved upon cooling (most noticeably for x=1and 0). Further isothermal TGA measurements at various temperatures havealso shown this kinetically oxygen-content limiting behavior, whichincreases equilibration time at lower temperatures (this limitingbehavior accounts for the significant differences in absorption rates ofFIGS. 4 and 7). Therefore, the OSC of samples (x=0 and 1) increase atlower isothermal temperatures and the desorption rate of x=0.7 increasesat slightly higher temperatures. The nature of these transitions fromthe P6₃ cm phase (δ=0) to the Hex₂ phase (δ=0.25) and from the Hex₂phase to the Hex₃ phase (δ=0.40) appears to easily equilibrate tointermediate oxygen content values. As a result, a mixture of severalphases will occur in various oxygen partial-pressures and temperatures,where low-temperatures, 150-200° C., can favor the Hex₃ phase;intermediate-temperatures, 230-330° C., favor the Hex₂ phase; andhigh-temperatures, above ˜275-375° C., favor the stoichiometric P6₃ cmphase (these ranges are dependent on oxygen partial-pressure and Dy/Ycontent). The slope of oxygen content versus temperature during the P6₃cm-Hex₂ phase transition at constant temperature (as well as on coolingin FIG. 4) decreases with increased Y content, which again indicatesslower absorption rates of Y-rich samples. Direct comparisons of theseabsorption rates are, however, complicated by slower oxygen ion kineticsat lower temperatures, which can approximated by

  D = D₀? ?indicates text missing or illegible when filed

The lower temperatures at which the Hex₂-P6₃ cm phase transition occursfor Y rich samples prevents absorption comparisons at similartemperatures; thus, the differences in absorption observed in FIG. 7 aredue to both differences in activation energy and temperature. Thisincreased rate of transition from the P6₃ cm phase to the Hex₂ phase mayalso be due to increased distortion to the P6₃ cm structure caused bylarger average R-site anions. On the other hand, the transition from theHex₂ to Hex_(a) phase (δ≧˜0.25) appears to favor Y doped DyMnO_(3.25)samples (x=0.1, 0.3, 0.5) over pure DyMnO_(3.25), as seen on cooling inFIG. 4.

Hydrogen reductions in TGA for DyMnO_(3+δ) and YMnO_(3+δ), which wereinitially done to determine oxygen content, showed to have increasedstability on reduction at δ=−0.12 and −0.20, respectively (as seen inFIG. 3 for DyMnO₃). To test for recoverability of the P6₃ cm DyMnO_(3+δ)and YMnO_(3+δ) phases, materials were heated to and held at 400° C. in42% H2/Ar in TGA until these respective values of δ were reached. Thesesamples were then cooled in Ar to 330 and 230° C., respectively, andheld at these temperatures under O₂. Samples quickly returned tostoichiometric oxygen content (>1 hour) and continued to absorb oxygenas seen during oxygen cycles in FIG. 7. XRD measurements after thisprocess confirmed that samples did not decompose to simple oxides. Thus,the addition of cycling to 400° C. in hydrogen to either thermal oroxygen partial-pressure cycling would yield an additional ˜450-1050μmol-O/g (for x=0-1) and would place these materials up to near recordlevels of OSC, ranging from 1150-2650 μmol-O/g (TABLE 2, wherecalculated values assume the stabilities seen at δ=−0.12 to δ=−0.20changes proportionally with x for intermediate samples).

While the values measured here do not surpass the best observed OSC inthe literature and the slow oxygen kinetics of Y-rich samples (x=0.7,0.9, 1) may be a limiting factor for their potential use for OSCapplication, the Dy_(1-x)Y_(x)MnO_(3+δ) system does have the several keyadvantages for application over these other candidates. First andforemost, the Dy_(1-x)Y_(x)MnO_(3+δ) system has the lowest reportedreduction temperature, being approximately 25-125° C. lower than therecord reduction temperature of YBaCO_(4-x)Al_(x)O_(7+δ) (withsignificant OSC values). On further comparison toYBaCO_(4-x)Al_(x)O_(7+δ), which decomposes at 550-700° C.,Dy_(1-x)Y_(x)MnO_(3+δ) has far superior stability, remaining stable upto 1100-1400° C. Additionally, from a hazardous waste and coststandpoint, mass-production of manganese oxides is much preferable tothat of cobalt or chromium oxides. Finally, there is great potential forthe Mn cation in hexagonal RMnO_(3+δ) to have large changes in oxidationstate because, unlike the majority of OSC materials, which depend on thecreation of oxygen ion vacancies or interstitial sites athigh-temperatures, the hexagonal Dy_(1-x)Y_(x)MnO_(3+δ) (as seen alsowith YBaCO_(4-x)Al_(x)O_(7+δ)) relies on reversible phase transitionsbetween several structures containing transition metal ions in variablecoordination. The potential OSC of related hexagonal manganites couldeasily surpass the current highest reported values, if they can bemodified to easily and reversibly transition in between phases withlarge amounts of Mn²⁺ and Mn⁴⁺ at low-temperatures.

Finally, apart from any possible OSC application, it should be notedthat hexagonal manganites have been largely believed to remainstoichiometric in oxygen content at elevated-temperatures. In situstructural measurements at high-temperatures have reported adisplacement of the MnO₅ bipyramids and a transition to the P6₃/mmcstructure, which occur for YMnO₃ at ˜650° C. and ˜950° C., respectively.Slight excesses of oxygen content (δ=0.01) have been reported at 1200°C. for YMnO_(3+δ) and ErMnO_(3+δ) but have not observed thenon-stoichiometric oxygen content behavior or the associated structuralchanges at lower temperatures as have been observed withthermogravimetric and XRD measurements. This behavior may not have beenpreviously observed in other hexagonal manganites due to the narrowrange of temperature (−200-350° C.) these new phases exist on heatingbefore returning back to δ=0 above ˜350° C. and the slow cooling or highoxygen partial-pressures they require. As discussed above, thistemperature range has not been of particular interest for structuralstudies of RMnO₃, as most of this work has been done at eitherlow-temperature to study magnetic ordering (≦200 K) or high-temperatureto measure the rattling behavior of the MnO₅ bipyramids or structuraltransitions (≧500° C.). The results herein indicate that the hexagonalRMnO_(3+δ) family is most likely prone to considerable oxygennon-stoichiometry and also show a direct relation between reductiontemperature and sorption rates of oxygen to the average ionic size of R.If this is the case, other hexagonal RMnO_(3+δ) materials withrare-earths that are close in ionic size to that of Y (e.g. Ho and Er)can have similar non-stoichiometric behavior. It should be noted thatthe synthesis of YMnO_(3+δ) under fast cooling to room temperatureyielded small, but measurable, excesses in oxygen content (δ=0.004).Many studies of RMnO_(3+δ) use samples prepared at elevated-temperaturefollowed by various cooling rates, which would yield slightlynon-stoichiometric samples for low-temperature measurements. Propertiesassociated with excess oxygen content (e.g. disruptions to the exchangeinteraction or the presence of Mn⁴⁺) may very well have had asignificant impact on the multiferroic properties of these samples, asit has been observed that even slight oxygen and cationnon-stoichiometry can have profound effects on magnetic and transportproperties of perovskite manganites.

Crystal Structure

XRD measurements were made to verify the hexagonal P6₃ cm structure ofDyMnO_(2.963) and DyMnO_(3.0) (samples 1 and 2) and to obtain apreliminary structural understanding of annealed samples. FIGS. 12A-12Fare a compilation of XRD patterns collected for samples 1*4, 6 and 7listed in TABLE 3. Peak positions and intensities of DyMnO_(2.963) andDyMnO_(3.0) were found to be in good agreement with previously reportedXRD patterns of P6₃ cm DyMnO₃. Furthermore, XRD data of the quenchedsample (sample 2) confirmed that stoichiometric samples are indeed thehexagonal P6₃ cm phase after quenching from above 400° C. as observedwith TGA data. XRD patterns of annealed samples (samples 3, 4, and 6) inthe 6 range of 0.18-0.24 clearly show growth of a second phase (Hex₂)and a disappearance of the P6₃ cm phase (where arrows indicate thegrowth and decrease of selected peaks for the P6₃ cm and the Hex₂ phase,respectively). The diffraction pattern of sample 6 (δ=0.24) is nearlysingle phase for this new set of peaks and is in agreement with thestability seen in TGA at δ ˜0.25 (FIG. 13). Finally, the XRD pattern ofthe high-pressure annealed sample (sample 7, δ=0.35) shows a decrease ofpeak intensity for the Hex₂ phase and the presence of additional peaks(third phase, Hex₃), which is again in agreement with TGA observations.The relative intensities of the Hex₂ and Hex₃ phases suggest that theHex₃ phase can have an oxygen content of δ=0.40, though this isdifficult to approximate due to the high degree of peak positionoverlapping. To help clarify the development of new peaks and peakoverlap, FIG. 14 shows an overlay of XRD patterns of samples 2, 6, and 7(δ=0.0, 0.24, and 0.35) in the 2θ range of 26-35°. FIGS. 13 and 14 showsimilarities of the diffraction patterns of the Hex₂, Hex₃, and P6₃ cmphases, which show that the Hex₂ and Hex₃ phases are structurallyrelated to the P6₃ cm phase. The increased number of peaks seen in theHex₂ and Hex₃ phases versus the P6₃ cm phase also shows a generallowering of symmetry or the formation of a super-structure. Finally, itshould also be noted, though these transformations are unlikely at theselow-temperatures under O₂, that the Hex₂ and Hex₃ phases were comparedto patterns of other known R_(x)Mn_(y) ⁴⁺Mn_(y-1) ³⁺O_(3+δ) systems(e.g. pyrochlore R₂Mn₂O₇, perovskite R-3c, R₂MnO₄ and RMn₂O₅ phases) andoxides (Mn₂O₃, MnO₂), which can account for the increase in oxygencontent. No traces of these structures were observed.

TABLE 3 List of annealed DyMnO_(3 + δ) samples. Conditions after Sampleno. synthesis of P6₃cm in Ar Sample type δ 1 None Small pellets −0.037 2Quenched from 420° C. air Small pellets 0.00 3 Cooled from 500° at Smallpellets 0.18 1.0° C./min in 21% O₂ at standard pressure 4 Cooled from500° at Small pellets 0.21 1.0° C./min in O₂ at standard pressure 5Cooled from 500° at Small pellets 0.24 0.1° C./min in O₂ at standardpressure 6 Cooled from 500° at Small pellets 0.35 0.1° C./min in O₂ at~250 bars

Guided by the initial XRD investigation, NPD measurements were conductedfor selected samples. High-resolution, backscattering data (2θ=144°,Bank 1 of SEPD) were used for DyMnO_(2.963), DyMnO_(3.0), andDyMnO_(3.21) (samples 1, 2, and 4, respectively) at room temperature.Low-angle scattering data (2θ=44°, bank 3) were also used forDyMnO_(3.21) at room temperature. High resolution synchrotron x-ray datawere also collected for DyMnO_(3.21) at room temperature.

Raw data for samples 1 and 2 were analyzed with the Rietveld method inthe space group P6₃ cm based on previous reports for the hexagonal RMnO₃system and the XRD measurements (FIG. 15). Structural sites of thisrefinement were 2a for Dy1 and O3; 4b for Dy2 and O4; and 6c for Mn1,O1, and O2. Cations occupancies were fixed at one and the siteoccupancies of oxygen ions were allowed to vary. Initial refinements ofthe DyMnO_(3.0) sample's occupancies varied less than one standarddeviation from fully stoichiometric oxygen content and were fixed to onefor its final fitting. For DyMnO_(2.963) occupancies of the O1 and O2sites were also fixed to one as refinements yielded values slightlygreater than one. Oxygen ion vacancies were found to prefer the O3 andO4 sites nearly equally. The resulting oxygen content of this samplecalculated from these refined occupancies (δ=−0.045) is in reasonableagreement with the value obtained from TGA (δ=−0.037). For both thesesamples (1 and 2), the calculated diffraction pattern of P6₃ cm is ingood match with the observed data for both samples and their latticeparameters are in agreement with a previous XRD report for DyMnO₃. Bondlengths were calculated using the geometric average by assuming fullsite occupancy. The average (Mn—O) bond length clearly increases fromthe stoichiometric to the reduced state, while the average (Dy—O) bondlength remains, relatively, unchanged. Again, this is due to theenlargement of the Mn^((3+2δ)+) cation with increasing oxygendeficiency. These results are in agreement with the oxygen vacancydependence of the tolerance factor and support the synthesis argumentsof forming the hexagonal phase by reduction of the perovskite RMnO_(3+δ)phase.

Analysis of neutron and synchrotron diffraction data for sample 4(DyMnO_(3.21)) revealed the formation of a large superstructureconstructed by tripling the c-axis of the P6₃ cm phase (c>33 Å). Severalother superstructure models and combinations of possible phase mixtureswere also examined but they all failed to index the large number ofextra peaks. Analysis of the superstructure's structural symmetry led tothe identification of R3 as the space group that could successfullyindex all peaks including the tiny ones. We note here that there is nodirect relationship between the two P6₃ cm and R3 space groups. Such agroup/subgroup relationship is not required for two samples that are notchemically the same. A group/subgroup relationship is required whendealing with a unique sample in which structural phase transitions occurat various temperatures or pressures. In the present case, the R3structure of the oxygen loaded DyMnO_(3.21) sample was determined as thespace group of the highest symmetry that can be successfully used toindex all Bragg reflections and refine the positions of the Dy and Mncations. Determination of the exact locations and site occupancies ofthe diverse oxygen atoms remain challenging due to the complexity of thesuperstructure and the nature of synchrotron x-rays that are inherentlymuch less sensitive to oxygen than neutrons, especially in the presenceof Mn and the heavy Dy rare-earth. Rietveld refinements usingsynchrotron data are presented in FIGS. 16A and 16B. In the refinements,the cation positions and thermal factors were all refined whereas theoxygen atoms were kept fixed at positions derived from the tripledstructure. As shown in FIGS. 16A and 16B, two phases were included inthe final refinements: the small parent P6₃ cm hexagonal structure(lower tick marks) and the larger R3 superstructure (upper tick marks).Fractional percentages by weight for the two phases refined to 14% and86%, respectively. It's obvious that the parent phase fails to index theobserved extra peaks that refine with R3. The superstructure's latticeparameters are listed in TABLE 4 together with the positions of the Dyand Mn cations in which we have high confidence. The exact determinationof the oxygen atoms in such a small molecule-like superstructure wouldnecessitate further collection of high quality neutron diffraction datapreferably using new RMnO_(3+d) samples in which the highly neutronabsorbing Dy would be replaced by Y or other trivalent rare-earthelements with significantly smaller neutron absorption cross sectionssuch as Ho and Er.

TABLE 4 Structural parameters for the R3 superstructure of DyMnO_(3.21).R3 DyMnO_(3.21) Atom x y z B (Å²) Dy 1 0 0 −0.06995(5) 030(4) Dy2 0 0 0.07051(5) 0.04(4) Dy3 0 0  0.25180(11) 0.19(3) Dy4 0 0  0.43022(6)0.97(5) Dy5 0 0  0.57090(6) 0.50(4) Dy6 0 0  0.74966(11) 0.32(3) Mn10.4261(5) 0.0019(9)  0  2.8(1)^(a) Mn2 0.3672(7) 0.6254(6)  0.50109(21)0.47(6)^(a) Lattice parameters (Å) a = 6.231(4) and c = 33.346(3)Reliability factors R_(wp) = 13.7%, R_(p) = 9.8%, R₁ = 4.4%, C² = 11.4^(a)These values clearly correlate with the undetermined distortedoxygen environment around Mn as expected from the insertion offractional amounts of additional oxygen. Please see the text for moredetails.

Thermal and Chemical Expansion

Expansion of the crystal lattice can occur through two mechanisms:thermal and chemical expansion. Thermal expansion (TE), as discussed intolerance factor arguments, is caused by expansion of the (R—O) and(Mn—O) bond lengths due to increased thermal energy at elevatedtemperature. Chemical expansion (CE) is caused by expansion of thelattice due to changes in oxygen stoichiometry. The TGA measurements ofDy_(1-x)Y_(x)MnO_(3+δ) materials, discussed above, have shown largechanges in oxygen stoichiometry between two stable oxygen contentregions, which occur on heating over a relatively short time scale (≦2hours) and narrow range of temperatures (˜100° C.). Thesecharacteristics allowed for the measurement of the effective CE over anarrow range of temperatures by simply subtracting the relatively smallvalue of TE from the observed value of CE. Similarly precisemeasurements of TE, without the any effect from CE, were possible intemperature regions of stable oxygen content. It should also be notedthat in some cases the thermal expansion coefficient (TEC) is consideredto be the net result of both CE and TE; here these are considered to beseparate effects, thus TEC herein is only attributed to TE. Thefollowing equations were used to calculate TE and CE:

${TEC} = {\frac{1}{L_{0}}\frac{1}{n - m}{\sum\limits_{i = m}^{n}\frac{{\Delta \; L_{i + 1}} - {\Delta \; L_{i}}}{T_{i + 1} - T_{i}}}}$

measured in K⁻¹, where L₀, ΔL, and T are the sample starting length, thechange in length, and temperature, respectively, and m-n are the setsfrom the measured temperature ranges and

${{CE} = {\frac{1}{\Delta \; \delta}\left( {{\frac{\Delta \; L}{L_{0}} -} < {TEC} > {\Delta \; T}} \right)}},$

measured in (moles of O)⁻¹, where Δδ is the absolute change in oxygencontent from stoichiometric 3.0 and <TEC> is the average TEC of the twooxygen content stable regions.

A perovskite sample of DyMnO₃ for dilatometry was cut from a densepellet after initial synthesis in air (˜5×3×2 mm in shape) and wasmeasured in 21% O₂/Ar atmosphere with heating rates of 0.5° C./min to900° C. (FIG. 11). Previous studies of the perovskite DyMnO_(3+δ) phasehave shown that it remains stoichiometric in 21% O₂/Ar up to ˜1000° C.,thus the expansion seen in FIG. 8 is solely due to TE. The TEC wasmeasured from 50-850° C. and was found to be 7.3*10⁻⁶ K⁻¹, which is ingood agreement with a previous report.

Pellets for dilatometry measurements were cut from dense samples (x=0,0.3, 0.5, 0.7, 1) after synthesis of the hexagonal material (˜5×3×2 mmin shape) and were then annealed at 400° C. with 0.1° C./min cooling inO₂. The oxygen contents of these dense samples were also measured withidentical conditions on TGA to determine the appropriate temperatureranges to separately extract TE and CE coefficients (the structuralphases present and after dilatometry measurements were also confirmedwith XRD). TE values were measured for these samples in their respectivetemperature regions of stable oxygen content observed in TGA forδ=0.22-0.29 (˜50-300° C.) and for δ=0 (˜600-850° C.). CE values weremeasured during the reduction between these stable oxygen contents overapproximate temperature gradient of ˜100° C. in the range of 240-390°C., where approximately 90% of the total oxygen reduction occurs. FIGS.9A-9B show these measurements for DyMnO₃ and illustrates how thecombination of dilatometry and TGA measurements was used in to determineTE and CE for all Dy_(1-x)Y_(x)MnO_(3+δ) samples. The lower startingoxygen contents after annealing in oxygen and the slower reduction ofdense pellets (as seen for DyMnO_(3+δ) in FIG. 9A) versus the smallchucks of material observed in FIG. 4 during TGA measurement are due tothe differences in the samples' density, surface area, and diffusiondistances. The TEC of the hexagonal phases in these two temperatureregions of stable oxygen content were found to be quite different,8.2-10.2*10⁻⁶ K⁻¹ (δ=0.22-0.29) and 2.1-5.6*10⁻⁶ K⁻¹ (δ=0), whichindicates the TEC of the stoichiometric Hex₂ (δ=0.25) and P6₃ cm phaseare approximately 8.4-11.6*10⁻⁶ K⁻¹ and 2.1-5.6*10⁻⁶ K⁻¹, respectively(FIG. 10 a). The values of chemical expansion during loss of oxygencontent are 0.82-3.48*10⁻² mol⁻¹ (FIG. 10B), which increasesignificantly with Dy content.

Previous reports of single crystal hexagonal RMnO₃ materials (R═Y, Ho,Sc, and Lu) have shown to have lattice parameters that linearly increasein-plane and decrease along c with increasing temperature. Thecontraction of the c-axis has also been shown to increase for larger Rions. Thus, the effect of substantial contraction of the c-axis isresponsible for the observed small change of volume of the unit cell andsignificantly lowers TE of the polycrystalline P6₃ cm material whencompared to their Hex₂ or perovskite phases (7.3*10⁻⁶ K⁻¹ and 6*10-⁶ K⁻¹for the perovskite phase of DyMnO₃ and YMnO₃, respectively). It is alsoin agreement with the decrease of the net TE with increased Dy contentfor P6₃ cm materials as seen in FIG. 10A. This tendency is, however,reversed for the Hex₂ phase, which shows to have increased TEC withincreased Dy content. Finally, an increased rate of contraction alongthe c-axis at the Curie temperature, ˜650° C., was reported previouslyfor YMnO₃ and HoMnO₃ in one study, but was also not observed in anotherreport. No anomalous behavior near this temperature was observed;however, this effect can be beyond the sensitivity range of thedilatometer for a polycrystalline sample, where anisotropic effects areaveraged out. On the other hand, if dense hexagonal RMnO₃ materials arealso prone to small non-stoichiometric behavior on heating, as seen herefor the temperature range of 400-600° C. (0<δ<0.015), this effect can bedue to the CE associated with the reduction of a slightly oxygenatedsample to stoichiometric oxygen content. The measurements herein showthe importance of understanding oxygen content behavior, as slightchanges in oxygen content can have similar effects on the net expansionas structural changes, which are not associated with changes in oxygencontent (e.g., the P6₃ cm to P6₃/mmc phase transition).

The CE during transition from the mixed state Hex₂/P6₃ cm (˜85-100%,δ=0.22-0.25) materials to nearly single phase P6₃ cm has a much largereffect on total expansion than TE. The primary cause of the CE seenduring the P6₃ cm/Hex₂ transition is due to the change in ionic radiusof the Mn^((3+2δ)+) cation. Finally, for comparison, the CE valuesreported here are of the same order of magnitude as the CE associatedwith the absorption and desorption of oxygen in perovskite LaMnO₃ orsimilar substituted perovskite manganites (˜2.4*10⁻² mol⁻¹ and ˜1-4*10⁻²mol⁻¹). However, the effect of CE for the hexagonal structure is muchmore prominent than in the perovskite phase, due to the larger change inoxygen content occurring over a much narrower temperature range.

Conclusions

The results and previous work with perovskite manganites show that theincreasingly stronger reducing conditions are needed to form hexagonalDy_(1-x)Y_(x)MnO_(3+δ) with decreasing x (for x≦0.7). Previous reportsof synthesis of the perovskite phase from the hexagonal phase withsmaller rare-earths (Ho, Er, and Y) under high-pressure, support theargument that transformations occur at specific values of the tolerancefactor due to the temperature, oxygen non-stoichiometry, andcompressibility dependence of the (R—O) and (Mn—O) bonds lengths.Hexagonal Dy_(1-x)Y_(x)MnO_(3+δ) materials were observed to reversiblyabsorb large amounts of oxygen at ˜200-300° C. and to sharply desorbthis uptake of oxygen during transition back to the stoichiometric P6₃cm phase above ˜275-375° C. or lower temperatures in lowerpartial-pressures of oxygen. Increased reversible changes in oxygencontent were achieved by annealing at high-pressures (δ=0.25-0.35) andwith hydrogen reduction at 400° C. (δ=−0.12-−0.20), which, if combined,can yield reversible oxygen storage capacities up to ˜2650 μmol-O/g.Rates of oxygen absorption were also observed to significantly decreasewith increasing yttrium content. The non-stoichiometric oxygen contentof these hexagonal manganites no doubt has profound influence on theirmultiferroic properties.

REFERENCES

-   Shelley, S. Chem. Eng. Prog. 2009, 105, 6.-   Ka{hacek over (s)}par, J.; Formasiero, P.; Hickey, N. Catal. Today    2003, 77, 419.-   Kodama, T.; Gokon, N. Chem. Rev. 2007, 107, 4048.-   Xu, Z.; Qi, Z.; Kaufman, A. Power Sources 2003, 115, 40.-   Sakakini, B.; Taufig-Yap, Y.; Waugh, K. J. Catal. 2000, 189, 253.-   Ciferno, J.; et al. DOE/NETL-2007/12912007.-   Rydén, M.; Lyngfelt, A.; Mattisson, T.; Chen, D.; Holmen, A.;    Bjørgum, E. I. J. Greenhouse Gas Control 2008, 2, 21.-   Klara, J.; et al. DOE/NETL-2008/13072007.-   Readman, J.; Olafsen, A.; Larring, Y.; Blom, R. Mater. Chem. 2005,    15, 1937.-   Figueroa, J.; Fout, T.; Plasynski, S.; McIlvried, H.; Srivastava, R.    I. J. of Greenhouse Gas Control 2008, 2, 9.-   Pei, S.; Kleefisch, M.; Kobylinski, T.; Faber, J.; Udovich, C.;    Zhang-McCoy, V.; Dabrowski, B.; Balachandran, U.; Mieville, R.;    Poeppel, R. Cata. Lett. 1995, 30, 201.-   He, H.; Dai, H. X.; Au, C. T. Catal. Today 2004, 90, 245.-   DiMonte, R.; Formasiero, P.; Graziani, M.; Ka{hacek over    (s)}par, J. J. Alloys and Comp. 1998, 275, 887.-   Nagai, Y; Yamamoto, T.; Tanaka, T.; Yoshida, S.; Nonaka, T.;    Okamoto, T.; Suda, A.; Sugiura, M. Catal. Today 2002, 74, 225.-   Singh, P.; Hegde, M.; Gopalakrishnan, J. Chem. Mater. 2008, 20,    7268.-   Karppinen, M.; Yamauchi, H.; Otani, S.; Fujita, T.; Motohashi, T.;    Huang, Y.; Valkeapää, M.; Fjellvag, H. Chem. Mater. 2006, 18, 490.-   Motohashi, T.; Kadota, S.; Fjellvag, H.; Karppinen, M.; Yamauchi, H.    Mater. Sci. Eng. B2008, 148, 196.-   Kadota, S.; Karppinen, M.; Motohashi, T.; Yamauchi, H. Chem. Mater.    2008, 20, 6378.-   Räsänen, S.; Motohashi, T.; Yamauchi, H; Karppinen, M. J. Solid    State Chem. 2010, 183, 692.-   Chmaissem, O.; Zhen, H.; Huq, A.; Stephens, P.; Mitchell, J. J.    Solid State Chem. 2008, 181, 664.-   Rydén M.; Lyngfelt, A.; Mattisson T.; Chen, D.; Holmen, A.;    Bjørgum, E. I. J. Greenhouse Gas Control 2008, 2, 21.-   Readman, J.; Olafsen, A.; Larring, Y.; Blom, R. J. Mater. Chem.    2005, 15, 1937.-   Motohashi, T.; Ueda, T.; Masubuchi, Y.; Takiguchi, M.; Setoyama, T.;    Oshima, K.; Kikkawa, S. Chem. Mater. 2010, 22, 3192.-   Yakel, H. L.; Koehler, W.; Bertaut, E.; Forrat, E. Acta. Cryst.    1962, 16, 957.-   Yakel, H. L. Acta. Cryst. 1955, 8, 394.-   Shannon, R. D. Acta. Cryst. A 1976, 32, 751.-   Yakel, H. L.; Koehler, W. C.; Bertaut, E. F.; Forrat, E. F. Acta.    Cryst. 1963, 16, 957.-   Dabrowski, B.; Chmaissem, O.; Mais, J.; Kolesnik, S.; Jorgensen, J.    D.; Short, S. J. Solid State Chem. 2003, 170, 154.-   Dabrowski, B.; Kolesnik, S.; Baszczuk, A.; Chmaissem, O.; Maxwell,    T; Mais, J. J. Solid State Chem. 2005, 178, 629.-   Kamegashira, N.; Satoh, H.; Ashizuka, S. Mater. Sci. Forum 2004,    449, 1045.-   Park, J.; Park, J. G.; Jeon, G. S; Choi, H. Y.; Lee, C.; Jo, W.;    Bewley, R.; McEwen, K. A.; Perring, T. G. Phys. Rev. B2003, 68,    104426.-   Lee, S.; Pirogov, A.; Han, J. H.; Park, J. G.; Hoshikawa, A.;    Kamiyama, T. Phys. Rev. B2005, 71, 180413(R).-   Ivanov, V. Y.; Mukhin, A. A.; Prokhorov, A. S.; Balbashov, A. M;    Iskhakova, L. D. Phys. Solid State 2006, 48, 1726.-   Carp, O.; Patron, L.; Ianculescu, A.; Pasuk, J.; Olar, R. J. Alloys    and Comp. 2003, 351, 314.-   Szabo, G.; Paris, R. A. Seances Academy Sci. C1969, 268, 517.-   Brinks, H. W.; Fjellvag, H.; Kjekshus, A. J. Solid State Chem. 1997,    129, 334.-   Suescun, L.; Dabrowski, B.; Mais, J.; Remsen, S.; Richardson Jr., J.    W.; Maxey, E. R.; Jorgensen, J. D. Chem. Mater. 2008, 4, 1636.-   Zhou, J. S.; Goodenough, J. B.; Gallardo-Amores, J. M.; Morán, E.;    Alario-Franco, M. A.; Caudillo, R. Phys. Rev. 82006, 74, 014422.-   Waintal, A. J. Chenavas 1967, 2, 819.-   Tachibana, M.; Shimoyama, T.; Kawaji, H.; Atake, T.;    Takayama-Muromachi, E. Phys. Rev. B2007, 75, 144425.-   Uusi-Esko, K.; Malm, J.; Imamura, N.; Yanauchi, H.; Karppinen, M.    Mat. Chem. Phys. 2008, 112, 1029.-   Lonkai, Th.; Tomuta, D. G.; Amann, U.; Ihringer, J.; Hendrikx, R. W.    A.; Többens, D. M.; Mydosh, J. A. Phys. Rev. B2004, 69, 134108.-   Jeong, I.; Hur, N.; Proffen, T. J. App. Cryst. 2007, 40, 730.-   Kamata, K.; Nakajima, T.; Nakamura, T. Mat. Res. Bull. 1979, 14,    1007.-   Katsufuji, T.; Masaki, M.; Machida, A.; Moritomo, M.; Kato, K.;    Nishibori, E.; Takata, M.; Sakata, M.; Ohoyama, K.; Kitazawa, K.;    Takagi, H. Phys. Rev. B2002, 66, 134434.-   Zhou, J. S.; Goodenough, J. B.; Gallardo-Amores, J. M.; Morán, E.;    Alario-Franco, M. A.; Caudillo, R. Phys. Rev. B2006, 74, 014422.-   Fiebig, M.; Lottermoser, T.; Pisarev, R. V. Appl. Phys. 2003, 93,    8194.-   Vajik, O. P.; Kenzelmann, M.; Lynn, J. W.; Kim, S. B.; Cheong, S. W.    Phys. Rev. Lett. 2005, 94, 087601.-   Lonkai, Th.; Tomuta, D. G.; Amann, U.; Ihringer, J.; Hendrikx, R. W.    A.; Többens, D. M.; Mydosh, J. A. Phys. Rev. B2004, 69, 134108.-   Jeong, I.; Hur, N.; Proffen, T. App. Cryst. 2007, 40, 730.-   Rao, C. N. R.; Serrao, C. R. Mater. Chem. 2007, 17, 4931.-   Choi, W. S.; Kim, D. G.; Seo, S. S. A.; Moon, S. J.; Lee, D.;    Lee, J. H.; Lee, H. S.; Cho, D. Y.; Lee, Y. S.; Murugavel, P.; Yu,    J.; Noh, T. W. Phys. Rev. B2008, 77, 045137.-   Nandi, S.; Kreyssig, A.; Yan, J.; Vannette, M.; Lang, J.; Tan, L.;    Kim, J.; Prozorov, R.; Lograsso, T.; McQueeny, R.; Goldman, A. Phys.    Rev. B2008, 78, 075118.-   Dabrowski, B.; Klamut, P. W.; Bukowski, Z.; Dybzinski, R.;    Siewenie, J. E. J. Solid State Chem. 1999, 144, 461.-   Bukowski, Z.; Dabrowski, B.; Mais, J.; Klamut, P. W.; Kolesnik, S.;    Chmaissem, O. J. App. Phys. 2000, 9, 5031.-   Zhou, H. D.; Denyszyn, J. C.; Goodenough, J. B. Phys. Rev. B2005,    72, 224401.-   Remsen, S. Ph.D. Dissertation, Northern Illinois University, 2010.-   Fu, B.; Huebner, W. Mater. Res. 1994, 9, 2645.-   Chen, X.; Yu, J.; Adler, S. B. Chem. Mater. 2005, 17, 4537.-   Miyoshi, S.; Hong, J.; Yashiro, K.; Kaimai, A.; Nigara, Y.;    Kawamura, K.; Kawada, T.; Mizusaki, J. Solid State Ionics, 2003,    161, 209.-   McIntosh, S.; Vente, J. F.; Haije, W. G.; Blank, D.; Bouwmeester, H.    Chem. Mater. 2006, 18, 2187.

1. A rare-earth manganese oxide, comprising M1, optionally M2, Mn and O,wherein: M1 is selected from the group consisting of In, Sc, Y, Dy, Ho,Er, Tm, Yb and Lu, M2 is different from M1, and M2 is selected from thegroup consisting of Bi, In, Sc, Y, La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy,Ho, Er, Tm, Yb and Lu, and Mn and O are present in an atomic ratio of1:z, and z is at least 3.15.
 2. The rare-earth manganese oxide of claim1, wherein z is at least 3.2.
 3. The rare-earth manganese oxide of claim1, wherein z is at least 3.25.
 4. The rare-earth manganese oxide ofclaim 1, wherein z is 3.15 to 3.4.
 5. The rare-earth manganese oxide ofclaim 1, wherein M1 and M2 are present in an atomic ratio of x:1−x, andx=0.1 to
 1. 6. The rare-earth manganese oxide of claim 5, wherein x=0.3to
 1. 7. The rare-earth manganese oxide of claim 1, wherein M1 isselected from the group consisting of Y and Ho.
 8. The rare-earthmanganese oxide of claim 1, wherein M1 is Y and M2 is selected from thegroup consisting of La, Ce, Pr, Nd, Sm, Eu, Gd, Tb and Dy.
 9. Therare-earth manganese oxide of claim 1, wherein M1 is Y and M2 is Dy. 10.The rare-earth manganese oxide of claim 5, wherein M1 and Mn are presentin an atomic ratio of 1:y, and 0<y≦10.
 11. A rare-earth manganese oxide,comprising M1, M2, Mn and O, wherein: M1 is selected from the groupconsisting of In, Sc, Y, Dy, Ho, Er, Tm, Yb and Lu, M2 is different fromM1, and M2 is selected from the group consisting of Bi, In, Sc, Y, La,Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb and Lu, M1 and M2 arepresent in an atomic ratio of x:1−x, and x=0.1 to 0.9, and Mn and O arepresent in an atomic ratio of 1:z, and z>3.
 12. The rare-earth manganeseoxide of claim 11, wherein z is at least 3.15.
 13. The rare-earthmanganese oxide of claim 12, wherein z is at least 3.2.
 14. Therare-earth manganese oxide of claim 12, wherein z is at least 3.25. 15.The rare-earth manganese oxide of claim 12, wherein z is 3.15 to 3.4.16. The rare-earth manganese oxide of claim 11, wherein x=0.3 to 0.9.17-20. (canceled)
 21. A rare-earth manganese oxide, comprising: (i) Mn,having a formal oxidation state between 3 and 4, (ii) O, and (iii) atleast one element selected from the group consisting of Bi, In, Sc, Y,La, Ce, Pr, Nd, Sm, Eu, Gd, Tb, Dy, Ho, Er, Tm, Yb and Lu, wherein therare-earth manganese oxide has an average temperature of maximum oxygenabsorption upon heating and cooling, T_(maxA), of at most 400° C., and atemperature of maximum oxygen desorption, T_(maxD), of at most 400° C.22-30. (canceled)
 31. An oxygen conducting membrane, comprising: (1) arare-earth manganese oxide, and (2) a support material, wherein themembrane has first and second opposing surfaces, the membrane is notpermeable to nitrogen gas, the rare-earth manganese oxide forms acontiguous structure exposed on both the first and second opposingsurfaces, and the rare-earth manganese oxide has an average temperatureof maximum oxygen absorption upon heating and cooling, T_(maxA), of atmost 400° C., and a temperature of maximum oxygen desorption, T_(maxD),of at most 400° C. 32-38. (canceled)
 39. An oxygen conducting membrane,comprising: (1) a manganese oxide, and (2) a support material, whereinthe membrane has first and second opposing surfaces, the membrane is notpermeable to nitrogen gas, the manganese oxide forms a contiguousstructure exposed on both the first and second opposing surfaces, andthe support material comprises at least one member selected from thegroup consisting of an organic polymer, a silicone rubber and glass.40-44. (canceled)
 45. A method of preparing oxygen, comprising:separating oxygen from a mixture of gases containing the oxygen, byconducting the oxygen through a manganese oxide, or absorbing andreleasing the oxygen from the manganese oxide, wherein the separating iscarried out at a temperature of at most 400° C. 46-59. (canceled)